The Mechanical Properties, Microstructure Analysis and Damage Behavior of AlMg7 Matrix Composites Reinforced with α-Al₂O₃ Particles

Abstract

This research investigated the influence of volume fraction (30 vol.% and 40 vol.%) and particle size α-Al₂O₃ on the physical and mechanical properties of AlMg7 composites manufactured by the squeeze casting technique. The aim of the study was to characterize the microstructure, hardness, density, tensile strength (σmax), compressive strength (σcmax), and impact strength, with a discussion of the mechanisms of destruction. The obtained materials exhibited very low porosity (below 2%), confirming the high efficiency of the ceramic preforms infiltration process. It was found that both hardness and tensile strength increase with decreasing size of the reinforcing particles. The highest growth in hardness at 113% was observed for the composite with 40 vol.% of F1200 particles, while the highest tensile strength, 341 MPa, was noted for materials with 30 vol.% of the same fraction of α-Al₂O₃ particles. In the case of compressive strength, the opposite relationship was observed, where an increase in volume fraction to 40% resulted in a significant rise in σcmax to 522 MPa. The tests also indicated that an increase in the proportion of the brittle ceramic phase radically reduces the impact strength of composites compared to the matrix, which is typical for composite materials with a metallic matrix. Microstructure analysis of the fractures revealed that the mechanism of destruction depends on the type of load and the size and proportion of particles, which is reflected in the transition from transcrystalline cracking to delamination at the phase boundary. The results confirm that the strengthening processes of composites depend on the effective transfer of stresses at the microscopic level.

1. Introduction

In recent years, aluminum–magnesium (Al-Mg) alloys have attracted considerable attention because of their advantageous characteristics, including low density, high mechanical strength, good weldability, and excellent corrosion resistance. These attributes make them well-suited for demanding applications in aerospace, marine engineering, shipbuilding, automotive manufacturing, the chemical sector, and military technology [1,2,3,4].

Magnesium plays a decisive role in enhancing the mechanical performance of aluminum alloys. Song et al. [5] demonstrated that each 1% increase in Mg content can raise the tensile strength by about 35 MPa. A similar trend was observed by Borowski et al. [6], who reported that raising the magnesium content from 3.5% to 5% increased tensile strength from 220 MPa to 290 MPa. Zha et al. [7], using the Equal Channel Angular Pressing (ECAP) method, achieved tensile strengths of nearly 600 MPa combined with ductility of approximately 4.5%, highlighting the potential of Al-Mg alloys for high-performance applications. While increasing Mg content slightly lowers the elastic modulus, it simultaneously improves impact resistance [8]. Burzić et al. [9] further noted that in Al-Mg alloys the energy required for crack propagation grows faster than that for crack initiation, resulting in a notable increase in overall impact energy—from roughly 9.1–10.2 J in AlMg3 to about 30.4–32.7 J in AlMg7. Only a few studies have investigated in detail the mechanisms of energy absorption under dynamic loading, particularly in relation to the Mg fraction in aluminum composite matrices. This issue is especially relevant for ballistic protection. Endramawan et al. [10] confirmed that higher Mg content, along with increased sintering temperature, improves hardness and bending strength in Al-Mg/SiC composites. A higher Mg concentration also enhances the wettability of reinforcement particles, which strengthens interfacial bonding between phases [11]. Overall, magnesium additions improve hardness, flexural strength, tensile strength, and impact resistance in Al-Mg alloys, supporting their potential use in advanced engineering fields, including ballistic applications. Nevertheless, despite promising results, the available literature on high-Mg Al-Mg alloys remains limited, and there is no clear agreement on the optimum magnesium content.

A critical factor in the development of Al-Mg matrix composites is the precise selection of the matrix alloy, the type and volume fraction of reinforcement, as well as the manufacturing route and processing parameters. These choices directly determine the resulting density, porosity, and mechanical strength of the composite material [1,12]. The processing method in particular dictates the final microstructure, porosity level, uniformity of reinforcement particle distribution, and the integrity of phase boundaries—all of which strongly influence the mechanical performance of the composite [13]. Among the available fabrication techniques, squeeze casting has been shown to produce Al-based composites with very low porosity, which is essential for both structural durability and ballistic performance. Taşkın et al. [14] demonstrated that for SiC reinforcement levels of 15% and 20%, the particles were uniformly distributed, and porosity values remained below 1.5%, which is remarkably low. Comparable outcomes were reported by Giri et al. [15] for Al5083-based composites, where a 14% increase in density was observed. In contrast, much of the earlier research has focused on stir casting, a more economical process but one that imposes certain restrictions on the maximum achievable reinforcement content. More recent studies [16] confirm that squeeze casting provides superior material characteristics, establishing it as one of the most effective techniques for manufacturing high-performance composites. The hardness and tensile strength of Al-Mg composites are also closely linked to the efficiency of the degassing process. Achieving high-quality infiltration requires careful control of processing conditions in relation to the thermophysical properties of the preform [17,18,19].

Furthermore, controlling the chemical composition of the aluminum matrix and the appropriate choice of reinforcement are essential for maximizing performance [20]. The addition of Al₂O₃ particles, for example, has been demonstrated to further enhance hardness and improve overall strength.

Vinodh et al. [21] reinforced AA5083 alloy with SiC and B₄C particles using the friction stir welding (FSW) technique. Their study demonstrated that the introduction of 6% reinforcement substantially enhanced tensile strength, hardness, and fatigue life. However, increasing the content to 8% led to greater brittleness and reduced impact resistance, underscoring the trade-off between strength and ductility. Similar observations were reported by Siddaraju et al. [22], who showed that adding up to 6% Al₂O₃ to AA5083 alloy improved tensile strength, yield strength, and microhardness, although it reduced elongation. Chandrashekar et al. [23] investigated AlMg4.5 alloys reinforced with 2–8 wt% of Al₂O₃ nanoparticles, noting the maximum tensile strength of 203.95 ± 0.15 MPa at a 6% addition. Tosun and Kurt [24] compared SiC and Al₂O₃ reinforcements in Al-Mg alloys, reporting that composites strengthened with SiC exhibited greater hardness than those with the same amount of Al₂O₃.

Overall, these findings indicate that an optimal reinforcement level exists, beyond which the material becomes brittle and loses impact strength. However, this threshold varies depending on the reinforcement type and particle size, and requires further systematic investigation [14,15,21,23]. While most studies focus on relatively low reinforcement contents, only a few have examined aluminum matrix composites with high reinforcement fractions. For instance, Kurzawa and Kaczmar [25] showed that incorporating 30–40 vol.% Al₂O₃ led to significant improvements in flexural strength, while reinforcement levels of 40–55 vol.% SiC resulted in superior hardness, better wear resistance, and reduced risk of interfacial cracking between matrix and particles [26,27]. These results highlight the need for more in-depth research into composites with higher reinforcement contents. The growing interest in using aluminum matrix composites for ballistic protection further emphasizes this direction [8,28,29,30]. In such systems, the metallic matrix contributes ductility and corrosion resistance, whereas the ceramic phase provides the high hardness and stiffness required to deform and fragment incoming projectiles. Although the potential of these materials has been widely recognized, their optimization for specific ballistic applications—particularly with regard to microstructural design and control of phase interfaces—remains an active research challenge [24,28]. One of the main difficulties lies in maintaining sufficient matrix plasticity when the content of alloying elements is increased: while higher alloying levels enhance hardness, they can also induce brittleness [31].

Recent research suggests several promising methods to mitigate this drawback in high-Mg alloys. Cryogenic treatment [32,33,34] and two-stage forming [35] have both been shown to produce beneficial microstructural modifications. Increased strength has been attributed to higher dislocation density, while improved ductility results from a more uniform distribution and fragmentation of secondary-phase particles, which promotes energy absorption and reduces stress concentration. Notably, Heyduk et al. [36] reported that AlMg7 alloys can achieve tensile strengths as high as 569 MPa. These advances, though encouraging, still require deeper exploration in the context of reinforced composites. Understanding the effect of high Mg content on impact resistance and on the integrity of matrix–reinforcement interfaces is particularly important, as these factors are decisive under dynamic loading conditions.

Aluminum-based composites reinforced with aluminum oxide (Al₂O₃) offer a combination of relatively low density and high mechanical strength and stiffness [37]. Depending on the size, quantity, and distribution of ceramic particles, it is possible to modify the material’s response to both static and dynamic loads. Ballistic tests have revealed promising protective properties in composites containing both fine and coarse ceramic particles [2,28,29,31,38,39]. In papers [8,22], researchers indicated that the density, modulus of elasticity, and strength of the material increase with increasing particle content but at the expense of ductility, which is a key compromise in armor design. Zayed et al. [37] showed that the average hardness increased by up to 30% with an increase in ceramic particle content. Also in work [40], it was demonstrated that adding nano Al₂O₃ particles to Al-Mg alloy results in an increase in hardness and compressive strength. Furthermore, research by Yu et al. [41] and Ref. [23] proves that Al₂O₃ particles have a beneficial effect on corrosion resistance. This property is particularly important in applications where the material is exposed to an aggressive environment, such as protective armor. The studies by Yogavardhan et al. [42] and Renreng et al. [43] reveal that the smaller the size of Al₂O₃ particles lead to the higher hardness and strength of the material. Fine particles provide a larger contact surface with the matrix and more effectively block dislocation movement, leading to an increase in hardness. This is confirmed in [42], where hardness values of 93.5 HB for 120 µm particles and 121.64 HB for 20 µm particles are reported. However, it is essential to note that under dynamic conditions, larger particles can dissipate impact energy more effectively [44]. This mutual exclusion of properties necessitates detailed, comprehensive research that will allow for the development of a material with an optimal property profile in both static and dynamic conditions.

This work focuses on the study of the properties of new metal–ceramic composites produced by the squeeze casting process through the infiltration of porous ceramic preforms made of α-Al2O3 particles with a matrix in the form of AlMg7 aluminum alloy, which is characterized by a notably high magnesium content, representing an innovative approach. The research presented so far has focused on alloys with a lower Mg content and manufactured using a different technique. Manufacturing composites with a high ceramic phase content presents a technological challenge and creates new opportunities in the design of materials with increased hardness and stiffness, which is essential for ballistic applications.

The present study investigates how the metallic matrix microstructure and the interaction between ceramic and metallic phases affect the quasi-static and impact behavior of Al-based composites. Key material properties—density, hardness, tensile and compressive strength, and impact resistance—are systematically measured to evaluate load transfer efficiency and the micromechanisms of failure. Emphasis is placed on correlating microstructural features with macroscopic performance to elucidate damage initiation and propagation under different loading regimes.

Composites containing high ceramic fractions (30 and 40 vol.% Al₂O₃) are produced and assessed as functional layers within multilayer ballistic concepts. In such assemblies, the ceramic-rich layer is intended to act as the primary impact element, promoting projectile deformation, fragmentation, and kinetic-energy dissipation.

To capture the influence of particle size on failure processes, both fine-grained and coarse-grained reinforcement distributions are compared, since earlier works report superior static properties for finer fractions but do not fully explain dynamic failure modes [42,43].

Experimental results will serve as calibration data for numerical models (e.g., Abaqus), enabling the development and validation of constitutive laws for ballistic simulations. Such model-based optimization aims to inform design choices for ceramic-metal multilayers by linking measured mechanical responses and observed microstructural damage to predictive simulations.

2. Materials and Methods

2.1. Materials

The choice of matrix alloy was guided by the target application of the composites as lightweight elements for ballistic protection. Accordingly, an AlMg7 alloy was selected as the matrix material. The relatively high magnesium content (7 wt%) provides not only excellent corrosion resistance but also markedly greater strength and hardness compared with conventional Al-Mg alloys such as AlMg4.5 or AW-5083. These enhanced properties are expected to assist in the initial deformation of incoming projectiles prior to the full engagement of the ceramic phase in energy absorption. The chemical composition of the AlMg7 alloy was verified using a GNR MiniLab 150 analyzer (Italy), and the obtained results are summarized in

Table 1. Chemical composition [wt%] of experimental alloys.

Element [% wt]	Mg	Mn	Si	Fe	Ti	Zn	Cu	Al
Results of analysis	7.15	0.92	0.06	0.10	0.02	0.32	0.02	balance

.

As the reinforcing phase, preforms composed of α-Al₂O₃ particles with size fractions F80, F150, F800, and F1200 were employed. The preforms contained 30 and 40 vol.% α-Al₂O₃, respectively, with the particle specifications, density, and size distribution summarized in

Table 2. Specification of selected α-Al₂O₃ particles for testing.

Chemical Composition [% wt]	α-Al2O3	SiO2	Fe2O3	Na2O	CaO	TiO2	K2O
	>99.0	<0.03	<0.04	<0.19	<0.01	<0.01	<0.01
Specific density: 3.95±0.5 [g/cm3]
Particle size: F80: 180 ÷ 212 [μm]; F150: 75 ÷ 106[μm]; F800: 6.5 ± 1% [μm]; F1200: 3 ± 1% [μm]

.

The morphology and EDS/EDX analysis results for selected α-Al₂O₃ particles (fractions: F80, F150, F800, and F1200) are shown in Figure 1. Observations confirmed their sharp-edged and irregular structure, typical of ceramic materials, after grinding.

2.2. Method of Composite Manufacturing

The composite materials were produced using the squeeze casting method from a liquid state to infiltrate porous ceramic preforms with an Al alloy [38]. In the first stage, preforms were made in the form of cylindrical plates with a diameter of 75 mm and a height of 14 mm. To ensure their geometric stability during the infiltration process, a hydrated sodium silicate solution (Na₂O⋅nSiO₂⋅xH₂O) was used as a binder to bond the ceramic elements. The preforms were pre-hardened by blowing with carbon dioxide (CO₂) and then stabilized by heating at a temperature of approximately 960 °C. The structure of the preform prepared in this manner, including the revealed binder bridges and the results of EDS/EDX analysis, is shown in Figure 2.

In the subsequent stage, the preforms were positioned within a stabilizing ring, which formed an integral part of the casting mold. After mold closure, liquid AlMg7 alloy was introduced and infiltrated the preforms under pressure applied by a punch. A pressing pressure of 90–100 MPa was maintained for approximately 30 s, until complete solidification of the casting. The process was conducted on a hydraulic press under the following conditions: liquid metal temperature of 720 °C, preform temperature of 480–500 °C, mold temperature of 320–350 °C, punch temperature of 120–150 °C, and pressing pressure of 90–100 MPa.

The diagram of the manufacturing stages for AlMg7-Al₂O₃ composites is shown in Figure 3.

Although Al-Mg alloys with Mg contents above 5 wt% are prone to segregation and the formation of brittle intermetallic phases, primarily β-Al₃Mg₂, the resulting microstructure displayed favorable features. This outcome is attributed to the squeeze casting process, which promotes rapid solidification under high pressure and leads to grain refinement in the metallic matrix. Consequently, a homogeneous structure was obtained, consisting of a fine-grained α-Al solid solution with β-Al₃Mg₂ precipitates dispersed throughout the matrix, occasionally forming branched morphologies. Within the infiltrated preform and in the interparticle regions, the β-Al₃Mg₂ phase was even more finely divided and uniformly distributed. Representative microstructures are presented in Figure 4.

2.3. Material Characterization Methods

To determine the density, the hydrophobic Archimedes method was used with toluene (density ρ = 0.87 g/cm³). The tests were conducted on an analytical balance at an ambient temperature of T = 20 °C. The results obtained are presented as volumetric density. In addition, the tests included the determination of porosity. To determine the experimental density of the materials, the classical method based on Equation (1) [43,45] was used:

$${\rho }_b={\displaystyle \frac{m_s}{m_s{-m}_{immersed}}}·{\rho }_c$$

(1)

where ${\rho }_{b }$—experimental density, $m_s$—mass of the dry sample in air [g], $m_{immersed }$—mass of the sample measured when completely immersed in the liquid, ${\rho }_c$—density of the liquid.

The porosity of the samples was determined in parallel using the method described in [43,45] and denoted according to Equation (2):

$$P={\displaystyle \frac{{\rho }_t-{\rho }_b}{{\rho }_t}}·100\%$$

(2)

where $P$—porosity, ${\rho }_{t }$—theoretical density, ${\rho }_b$—experimental density

The research was conducted at an ambient temperature of T = 20 °C using an analytical balance on rectangular samples with dimensions: l = 10–12 mm, w = 6 mm, h = 4 mm. Before analysis, the samples were dried at a temperature of T = 105 °C for 24 h until a constant mass $m_s$ was achieved.

Hardness tests were performed using the Brinell method (tester IM-700M Innovatest.

Europe BV, the Netherlands), using a ball with a diameter of 2.5 mm under a load of 625 N, in accordance with the PN-EN ISO 6506-1:2023-09 standard [46]. (Metals–Brinell hardness test–Part 1: Test method) [46].

Impact strength was determined using the Charpy method (VEB Werkstoffprüemaschinen, Leipzig), in accordance with PN-EN ISO 148-1:2017-02 [47]. Based on previous studies [48] on MMC composites with a high α-Al₂O₃ particle content, the samples were made without a notch, in the shape of a rectangular prism with the following dimensions: length l = 50 mm, height h = 4 mm, and width b = 6 mm. The samples’ dimensions were chosen to fit the available impact testing equipment, allow for repeatable preparation from a limited amount of material, and correspond to the castings from which they were cut.

Tensile strength tests were conducted on an INSTRON 3369 (Instron, High Wycombe, UK) machine at a speed of 2 mm/min [6]. The tests used samples with a nominal diameter of d = 6.0 mm, and a gauge length of l0 = 30.0 mm. Compression strength tests were conducted on cylindrical samples with a diameter of d = 6 mm, and a height of h = 9 mm. The dimensions of the samples and the strain rate were adopted in accordance with the requirements of PN-EN ISO 6892-1:2020-05 [49] and ASTM E9-19 [50], as well as the values used in studies in the literature [51]. All mechanical property tests were performed at room temperature, which was approximately 22−23 °C.

Observations and microstructure analyses of the materials were performed using a NIKON Eclipse MA200 (Nikon, Japan) light microscope coupled with an SSD digital camera and a JEOL JSM-IT210 (Jeol, Japan) scanning electron microscope SEM equipped with an EDS/EDX system. The studies were conducted on both metallographic cross-sections prepared according to ASTM E3 [52] and on fracture surfaces analyzed in an SEM. In the SEM analysis, a BSE detector was used in HV (High Vacuum) mode at an accelerating voltage of 15 kV and magnifications ranging from ×40 to ×15,000. The surfaces of the metallographic sections were prepared according to a procedure that included grinding on aqueous abrasive papers and polishing on polishing cloths using fine-grained suspensions. For better contrast, the sample surfaces were coated with graphite using a Q150T vacuum sputter system.

The matrix grain size was measured using the Nikon NIS Elements image analysis system (version 4.10).

3. Results

3.1. Identification and Microstructure

Microstructural characterization was carried out on polished metallographic cross-sections (Figure 5). The analyses confirmed that in AlMg-F80 and AlMg-F150 composites, a uniform particle distribution was achieved for both 30 and 40 vol.% Al₂O₃. In contrast, materials reinforced with finer F800 and F1200 fractions exhibited regions of unreinforced matrix circled in Figure 5c, with average dimensions of 6–10 μm, appearing between reinforced zones.

For the purpose of estimating the Hall–Petch strengthening effect, the grain size of the AlMg7 matrix material was manually measured. For this purpose, approximately 10 measurements were taken on 5 images—Figure 6. This way, the average grain size was determined, which was 17.74 μm.

Structural analysis of the composites was complemented by EDS/EDX mapping (Figure 7), which illustrates the elemental distribution on polished sample surfaces. In addition to the dominant matrix and reinforcement elements (Al, Mg, O), signals from silicon (Si) and sodium (Na) were also detected, originating from the silica-based binder used in preform fabrication.

Both the binder bridges connecting the particles and the silica coatings on their surfaces can react with constituents of the metallic matrix. Under the elevated temperature and pressure conditions of squeeze casting, even short exposure times promote atomic diffusion, resulting in the formation of intermetallic phases at particle–matrix interfaces. Figure 8 presents a line-scan analysis of elemental distribution, highlighting the presence of such phases as well as a developed binder bridge linking adjacent particles.

The presence of magnesium (Mg) in the alloy and silicon (Si) coming from silica bridges connecting ceramic particles, combined with a high infiltration temperature of 720 °C, creates conditions very favorable for the formation of a glassy silicate phase in the form of magnesium silicide (Mg₂Si) according to Formula (3) [53,54].

4Mg+SiO₂ → 2MgO+Mg₂Si

(3)

In addition, analyses have also revealed locally increased oxygen levels, suggesting the presence of magnesium oxide (MgO) in the structure, which may be formed as a result of magnesium reacting with trace amounts of oxygen in the matrix or at phase boundaries. It usually occurs in the form of small precipitations or thin layers and is formed by reaction with the matrix and particles according to (4) and (5) [11].

3Mg+Al₂O₃ → 3MgO+2Al

(4)

2Mg+O₂→2MgO

(5)

In the papers [53,55,56,57], researchers also confirmed the occurrence of these reaction products.

3.2. Density and Porosity

The incorporation of α−Al₂O₃ particles into the AlMg7 alloy matrix significantly increased the theoretical density of the material. For the base AlMg7 alloy, this density was 2.68 g/cm³, while for composites with 30 vol.% and 40 vol.% α−Al₂O₃, it increased to 3.06 g/cm³ and 3.19 g/cm³, respectively. This increase is consistent with the mixture rule (6) [58,59] and results directly from the higher density of ceramics compared to the metallic matrix.

$${\rho }_c=\sum \left(f_i{\rho }_i\right)=f_1{\rho }_1+f_2{\rho }_2+\cdots +f_n{\rho }_n$$

(6)

where ${\rho }_c$—density of the composite, ${\rho }_1, {\rho }_2, \dots ,{\rho }_n$—densities of each constituent in the composite, $f_1,f_2,\dots ,f_n$—volume fractions of each constituent

The AlMg7 base alloy exhibited the highest relative density (99.55%) and the lowest porosity (0.45%). In the composites, porosity values were slightly higher. For both reinforcement levels (30 and 40 vol.%), the lowest porosity was recorded for the F150 particle fraction, amounting to 0.64% and 0.79%, respectively (

Table 3. The influence of alloy strengthening with α-Al₂O₃ ceramic particles on density and porosity.

Materials	Theoretical Density [g/cm3]	Experimental Density [g/cm3]	Relative Density [%]	Porosity [%]
AlMg7	2.68	2.6680	99.55	0.45
F80-30 vol.% Al2O3	3.06	3.0095	98.32	1.68
F150-30 vol.% Al2O3	3.06	3.0413	99.36	0.64
F800-30 vol.% Al2O3	3.06	3.0359	99.18	0.82
F1200-30 vol.% Al2O3	3.06	3.0327	99.07	0.93
F80-40 vol.% Al2O3	3.19	3.1318	98.24	1.76
F150-40 vol.% Al2O3	3.19	3.1629	99.21	0.79
F800-40 vol.% Al2O3	3.19	3.1566	99.02	0.98
F1200-40 vol.% Al2O3	3.19	3.1550	98.96	1.04

Density particles Al₂O₃-3.95 [g/cm³].

, Figure 9a). This suggests that F150 particles enabled the most effective infiltration of the liquid alloy, thereby minimizing residual porosity. Interestingly, the results for F80 composites deviated from the general trend observed in the literature, where larger particles typically improve infiltration and reduce porosity. In this case, porosity values reached 1.68% (30 vol.%) and 1.76% (40 vol.%). Microstructural examination revealed that the F80 particles exhibited internal porosity, and in composites with 30 vol.% reinforcement, highly developed SiO₂ bridges were also present. The combination of porous particles and excessive binder bridges ultimately reduced the overall quality of these composites.

The porosity results obtained for the F800 and F1200 fractions, ranging from 0.82% to 1.04%, are consistent with general tendencies. They indicate that although smaller particles tend to form agglomerates that obstruct infiltration, the overall porosity of the material remains low. Comparing the results of reinforced materials for the F1200 with 30 vol.% the porosity is 0.93% and with 40 vol.% the porosity is 1.04%. It can be seen that increasing the ceramic content slightly increases the porosity. Ultimately, for all manufactured materials, the obtained porosity values were significantly below 2%, within the range that guarantees the preservation of good mechanical and physical properties of the manufactured MMC composites—Figure 9b.

3.3. Hardness

Figure 10 illustrates a marked increase in the hardness of the composites relative to the AlMg7 base alloy, which exhibited 91 HBW after solidification under pressure. This improvement is attributed to the incorporation of the hard ceramic phase. A clear trend was observed: hardness increased with decreasing reinforcement particle size, consistent with theoretical predictions that smaller particles provide a larger interfacial area and more uniform distribution, thereby enhancing strengthening [23,60,61]. For example, composites reinforced with F80 particles reached 124 HBW (30 vol.%) and 145 HBW (40 vol.%), whereas those with the finest F1200 fraction attained 181 HBW and 194 HBW, respectively. The hardness of the 40 vol.% F1200 composite was more than twice that of the AlMg7 matrix, corresponding to an increase of approximately 113%.

A second important trend concerns the effect of ceramic content. For all particle fractions, increasing reinforcement from 30 to 40 vol.% resulted in a further rise in hardness. For instance, F150 composites increased from 148 HBW to 156 HBW (an increase of approximately 5.4%), while F800 composites increased from 178 HBW to 188 HBW (an increase of approximately 5.6%). However, for the finest fractions F800 and F1200, the effect of further particle size reduction on hardness increase was less pronounced, as shown by the small difference between 30 vol.% F800 (178 HBW) and 30 vol.% F1200 (181 HBW), or between 40 vol.% F800 (188 HBW) and 40 vol.% F1200 (194 HBW). This indicates that below a certain particle size threshold, the efficiency of reinforcement by size reduction diminishes.

These observations are consistent with established strengthening mechanisms in metal matrix composites, including dispersion strengthening by fine, uniformly distributed particles and dislocation barrier effects, where particles impede dislocation motion [15]. The maximum hardness of 194 HBW, achieved for the composite containing 40 vol.% of F1200 particles, confirms that the combined effect of high ceramic content and fine particle size yields the greatest hardening effect.

3.4. Tensile and Compressive Test

The mechanical properties of Al alloys change radically after they are reinforced with ceramic particles. Alloys without ceramic content exhibit elastic-plastic properties, characterized by distinct flow and the possibility of large deformations [35,62]. In contrast, composites with an aluminum matrix reinforced with ceramic particles exhibit the characteristics of brittle materials, with significantly limited plasticity. In the materials studied, a 30% and 40% volume fraction leads to a dense distribution of particles, which effectively limits dislocation movement and the development of local plasticity zones, acting as mechanical barriers and load transfer initiators [11,15]. For this reason, their behavior under load is fundamentally different. In order to standardize the analysis and enable comparison of both groups of materials, the maximum compressive stress at which the samples are destroyed, and the maximum tensile stress at which they crack were adopted as common parameters. This approach allows for a uniform assessment of the strength of both ductile alloys and more brittle composites, providing a complete picture of their behavior under various load conditions. Due to the fact that unreinforced samples swell during compression without a clear maximum in the table, the test results are presented as a conventional yield stress Rp_0.2 (

Table 4. The influence of strengthening AlMg7 alloy with α-Al₂O₃ ceramic particles on the maximum tensile and compressive strength.

Materials	Maximum Tensile Strength σmax [MPa]			Maximum Compressive Strength σc max [MPa]
	min	max	Mean	min	max	Mean
AlMg7	178	203	193	152 *	165 *	161 *
F80-30 vol.% Al2O3	133	168	147	359	310	293
F150-30 vol.% Al2O3	160	221	203	383	431	395
F800-30 vol.% Al2O3	319	347	332	448	474	467
F1200-30 vol.% Al2O3	336	351	341	465	489	478
F80-40 vol.% Al2O3	-	-	-	314	380	373
F150-40 vol.% Al2O3	184	202	195	387	421	406
F800-40 vol.% Al2O3	248	261	254	476	527	514
F1200-40 vol.% Al2O3	275	296	289	509	533	522

*-yield strength (0.2% offset) Rp₀.₂.

).

The tests showed clear correlations in the behavior of materials between the type of loading, content, and size of Al₂O₃ reinforcing particles. The highest reinforcement efficiency, determined in relation to the unreinforced AlMg7 matrix with a base tensile strength σ_max = 193 MPa, was observed for composites with the finest Al₂O₃ particles, fractions F800 and F1200.

The tensile strength tests confirmed that the strengthening effectiveness of the composites is inversely proportional to particle size. The highest increase in strength was obtained for composites with the finest particles, F800 and F1200. For the composite with 30 vol.% of F800 particles, σ_max was 332 MPa, an increase of approx. 72% compared to the base alloy, while for the F1200 fraction, σ_max was 341 MPa, an increase of approx. 77%. This slight increase of 9 MPa for the finer fraction confirms the crucial role of the interfacial area in load transfer [11,63]. In the case of composites with 40 vol.%. particle content, a decrease in tensile strength is observed compared to composites with 30% reinforcement. For the F800 fraction, Rm 254 MPa, an increase of 32% was obtained, while for the F1200 fraction, σ_max was 289 MPa, a rise of 50%. The strength values for the F80 and F150 fractions decreased to a level equal to or lower than that of the unreinforced matrix, which indicates a loss of reinforcement effectiveness by large particles-Figure 11a. For the 40 vol.% Al₂O₃ F80 samples, no result was obtained due to the samples breaking in the holder.

In the case of compressive strength σc_max, the tests showed an inverse relationship with respect to σ_max between the amount of particles and the increase in stress transfer resistance—Figure 11b,c. In this case, the maximum strength value of 522 MPa was achieved by samples with the highest 40 vol.% of F1200 fraction particles, and 514 MPa for composites with the F800 fraction. Thus, in the particle size range of 3 ± 1% μm (F1200) and 6.5 ± 1% μm (F80), these results differed on average by only 8 MPa in favor of composites reinforced with the finer fraction. A noticeable difference in compressive strength compared to composites with 40 vol.% reinforcement was found in tests on composites with 30 vol.% reinforcement. In this case, the average value is lower by 44 MPa for F1200 composites and 38 MPa for composites with F800 particles. Tests on composites with F80 and F150 fractions showed a drastic decrease in strength compared to the strongest composites made with F1200 particles. For F150, this decrease is 18–23% for composites reinforced with 30 vol.% at 395 MPa and 478 MPa for composites with 40% reinforcement, respectively. The lowest strength value of 293 MPa was obtained for samples of 30 vol.% Al₂O₃ of the F80 fraction-Figure 11c.

Considering the specificity of the squeeze casting process and the solidification of the casting under high pressure (approx. 100 MPa), a reduction in the matrix grain size occurs. The strengthening contribution to the material’s yield strength resulting from this grain refinement effect, ${\Delta \sigma }_{H-P}$, can be quantified using the Hall–Petch relationship (7) [64,65].

$${∆\sigma }_{\mathrm{H}-\mathrm{P}}=k_y{d}^{-1/2}$$

(7)

where $k_y$—Hall–Petch coefficient; d—average grain size.

This strengthening mechanism, based on impeding the movement of dislocations by grain boundaries, is crucial for characterizing the AlMg7 matrix, whose average grain size was 17.7 μm. Based on this value and assuming $k_y$ = 0.136 MPa m^1/2 was calculated to be ${∆\sigma }_{\mathrm{H}-\mathrm{P}}=30.87 \mathrm{M}\mathrm{P}\mathrm{a}$.

In the composites, unlike in the unreinforced matrix, the solidification of the alloy occurs inside the preform in narrow, confined spaces between densely packed ceramic particles. The combination of solidification conditions resulting from the geometric structure of the preforms and process conditions, such as a high cooling rate, leads to an increase in the local solidification velocity. This, in turn, contributes to a further reduction in the matrix grain size within the reinforcement. It should therefore be assumed that the value of the ${∆\sigma }_{\mathrm{H}-\mathrm{P}}$ contribution originating from the matrix will further increase in the composite materials.

It should be emphasized that with respect to the CYS (Compressive Yield Strength) value of 161 MPa for the unreinforced matrix material, the dominant part of the matrix yield strength (approx. 130 MPa) is the result of solid-solution strengthening due to the high 7% Mg content.

Furthermore, the high CYS values are partially a result of the strengthening due to load transfer Δσ_L, which can be expressed by the simplified relationship (8) [64].

$$∆{\sigma }_L=0.5{\sigma }_m·V_p$$

(8)

where ${\sigma }_m$—yield strength of AlMg7, $V_p$—volume fraction of Al₂O₃

Although the simplified Equation is often derived in the context of tension, this mechanism is analogously effective under conditions of uniaxial compression [61,64].

The increase in Δσ_L directly correlates with the observed higher Compressive Yield Strength (CYS) of 472 MPa for the 40 vol.% particle fraction. Based on this relationship, the Δσ_L values obtained for the composites under investigation were 24.1 MPa for the 30 vol.% fraction and 32.2 MPa for the 40 vol.% fraction. The 8.1 MPa increase in the Δσ_L contribution, associated with increasing the particle content from 30 vol.% to 40 vol.%, is therefore a significant factor determining the higher strength of the composite with greater particle packing.

Furthermore, it has been demonstrated that smaller particles, such as F1200, improve strengthening and the compressive yield strength CYS more effectively than larger-sized particles, such as F80 or F150. This results directly from the smaller average particle diameter d, which leads to a larger contact area with the matrix material and presents a greater obstacle to dislocation movement, consistent with the strengthening phenomenon resulting from the Orowan effect [60,61,64,66,67,68].

It should be emphasized that overall strengthening is the result of the synergy of multiple components that must be taken into account.

Additionally, within the overall balance, particle boundaries and the significant difference in the coefficient of thermal expansion between the matrix materials and ceramic particles can generate regions with substantial stress fields. Fine precipitates of the β-Al₃Mg₂ phase form additional barriers to dislocation movement inside the grains. The combination of this effect with the presence of a high density of thermal dislocations caused by the difference in the coefficient of thermal expansion Δα, which is approximately 23 × 10⁻⁶ K⁻¹ for AlMg7 and 7.5 × 10⁻⁶ K⁻¹ for Al₂O_3, results in a significant increase in stress concentration at the phase interface and in the matrix, which further affects the mechanical and thermal properties of the composite.

The total strength of the composite is the result of the additive effect of various mechanisms, including those beyond the scope of this work. The increase in yield strength is the result of the combined contributions of ${\Delta \sigma }_{H-P}$ and ${\Delta \sigma }_L$, as well as the described Orowan strengthening effect and ${\Delta \sigma }_{CTE}$, the mentioned solid solution strengthening, and precipitation strengthening [61,64,67,68]. To obtain a complete and quantitative characterization of all strengthening components and their potential interactions, further research will be conducted.

Below, Figure 12 and Figure 13 show cross-sectional views of the samples in a parallel plane to the tensile force, at the point where the fracture metal was formed after the tensile test. Microscopic analyses of the samples subjected to σ testing revealed the influence of load transfer and dislocation pinning mechanisms, as well as the active participation of the particle–matrix interface in stress transfer processes under tensile loading. Literature research indicates that dislocation accumulation can contribute to an increase in the tensile strength and elongation of a material [66,67]. However, a high concentration of particles, especially in the area of phase boundaries, can lead to local stress concentrations, resulting in delamination at the matrix–particle interface and crack initiation, which corresponds to the observations presented in Figure 12, Figure 13 and Figure 14.

In all cases examined, the impact of tensile stresses on the effective cross-section is evident. The stress distribution within the material includes not only the fracture boundary but also regions deeper inside. Observations indicate that in materials reinforced with the largest particles, F80 (180–212 µm), the highest level of delamination occurs at the interphase boundary. Large particles anchored in the metallic matrix significantly restrict the material’s plastic deformation under tensile forces. This effect is amplified by the high packing density of the particles in the material (40% by volume), which significantly reduces the interparticle distances. Consequently, the matrix separates from the particles along the load direction, and the fracture follows the path of the weakest resistance line, choosing the interfacial boundary. In composites with F150 particles, there is notably better adhesion at the interfacial boundary, resulting in an increased proportion of transcrystalline cracking. This is evident from the visible fragmentation of the particles. The difference in fracture mechanisms can be explained by the reduction in the size of the binder bridges between particles in composites made from F150 particle preforms, compared to those with F80 particles. In this case, the smaller bridge size eliminates weak points at the phase boundary.

In materials containing smaller particles, such as fractions F800 and F1200, as in materials with F150, cracks propagate not only along interfacial boundaries but also transcrystalline through the particles themselves. This is evidenced by fragments of particles with sharp ends visible on the crack surface, which are still attached to the matrix. As a result of local stress concentrations in the area of the sample rupture, slight fragmentation occurs, visible on the surface of the fracture (Figure 13). Due to its specific microstructure, in which clusters of particles occur, the crack often propagates through the boundary between the cluster and the matrix or through the cluster itself.

The SEM tests and analyses carried out at ×50 magnification (Figure 14a,b) on the fracture surfaces of samples obtained after tensile strength testing revealed a visible significant change in fracture topography for materials with different particle sizes. In samples containing Al₂O₃ particles of the F80 and F150 fractions, the fracture surfaces show numerous high faults visible in the surface topography, associated with the passage of the fracture front through areas of local microdeformation. On the fracture surfaces of samples with F800 and F1200 particles, the faults are low, and the fractures are clearly flatter. The high volume of ceramic elements at 30% and 40 vol.% and their dense packing consequently led to a significant reduction in grain size, which resulted in a reduction in the size of the AlMg7 matrix material stretching. With an increase in the volume of reinforcing particles in the materials, the fracture surface becomes flatter with significantly lower faults.

In micro-areas without particles in composites with F800 and F1200 fractions, as shown in Figure 13c, the fracture surface changes its characteristics, becoming more irregular. The increased material stretching observed in these areas is the result of the larger grain size of the matrix compared to areas filled with particles (Figure 14d). In areas with dense particle packing, the fracture shows much greater regularity, which is characteristic of composites where crack propagation is hindered.

The convergent local distribution of Mg and Si confirms the presence of Mg₂Si on the surface of the fractures obtained after tensile strength (σ) testing (Figure 15).

It is worth noting that Mg₂Si actively participates in cracking mechanisms [53,55]. In this case, its presence in the form of lamellar cracked precipitations was found at the tops of the matrix material stretching that formed on the surface of the fractures (Figure 16). Additionally, sodium (Na) originating from silica bridges was observed in the analyzed fragment of the matrix. Sodium, present in trace amounts in this location, is dispersed in the matrix (

Table 5. Chemical composition analysis from the locations shown in Figure 16 (wt%).

Name	O		Al		Mg		Si		Na
	wt%	at.%	wt%	at.%	wt%	at.%	wt%	at.%	wt%	at.%
Spc_01	18.25	27.17	75.28	66.45	5.70	5.59			0.77	0.79
Spc_02	4.32	6.95	50.91	48.53	24.70	26.13	20.07	18.38
Spc_03	53.13	65.65	46.87	34.35

).

During the compression of unreinforced AlMg7 alloy samples, permanent swelling was observed without the appearance of macroscopic cracks. Whereas in the composites reinforced with ceramic particles, destruction occurred through shear cracks oriented at an angle of approximately 45° to the compression axis (Figure 17). This type of fracture is consistent with the theory of destruction mechanisms, where cracks are initiated in planes of maximum shear stress [69].

Composites with 30 vol.% of particles exhibited greater plastic deformation capacity compared to samples containing 40 vol.% of particles. This phenomenon is the result of a lower proportion of particles in the metallic matrix, which allows for the freer movement of dislocations, which is crucial for plastic properties. In these areas, the particles constituted a clear obstacle to deformation. This was evidenced by extensive deformation zones in the vicinity of the main fracture. Within these zones, numerous particles effectively transferred stresses and then fractured. Matter transfer was also observed in these zones, resulting in displacements and delaminations along previously fractured particles, even at greater distances from the main fracture (Figure 17c–g).

In some samples with 30% particle volume, no clear fracture occurred, and a distinct plastic deformation band was observed where particle transfer and rearrangement took place, as illustrated in Figure 17c, for example. Numerous particle fractures were observed in this area.

In composites manufactured on preforms with smaller particles, F800 and F1200, microcracks formed within the main crack, running approximately parallel to the central fracture plane. In these zones, the deformations propagated along preferred boundaries, separating micro-areas not reinforced with particles from areas with particles and through particle clusters. The studies also revealed cracked Mg₂Si phase precipitates in the plastic deformation zone (Figure 17f).

The observations reveal that in samples tested for compressive strength (σc), as in the case of tensile strength (σ), the influence of load transfer and dislocation pinning mechanisms and inter-particle zone stiffening was significant. The high proportion of the reinforcing phase resulted in a significant reduction in plastic deformation capacity, which in the samples with the highest degree of strengthening led to the initiation and propagation of cracks.

3.5. Impact Strength

As part of the impact testing conducted using the Charpy method, the energy absorption work required to initiate and propagate a fracture was determined. It was found that a high ceramic phase content had the opposite effect to that observed for hardness HBW and compressive strength σc–it significantly reduced the composite’s ability to absorb energy under dynamic loads. All tested composite samples indicated substantially lower resistance to dynamic loads compared to the material without particles, AlMg7 alloy, whose impact strength was 72.7 J/cm². At the same time, the tests confirmed the significant influence of both the quantity and size of the reinforcing particles on impact strength, as shown in Figure 18. The influence of the Al₂O₃ reinforcement content and particle size in the AlMg7 alloy matrix on the total impact energy is shown in Figure 18a, while the impact strength, defined as the ratio of the work used to destroy the sample to its cross-sectional area, is shown in Figure 18b.

It was found that the total energy absorbed by the composites decreases with an increasing volume fraction of Al₂O₃ particles and increases with a decreasing particle size.

The AlMg7 matrix alloy demonstrated the highest total energy value for destruction, achieving an average of 17.3 J. The addition of 30 vol.% of Al₂O₃ particles with sizes F800 and F1200 resulted in a decrease in this value by approximately half to 8.5 J and 8.8 J, respectively. In the case of composites containing F150 particles, the total energy decreased by approximately three times, whereas the use of F80 particles resulted in an Et value dropping to about 2.1 J, which corresponds to an eightfold reduction compared to the unreinforced matrix.

Introducing 40 vol.% Al₂O₃ particles, similar relationships were obtained, with the total energy values being approximately 5% lower for F150 and 15% lower for F800 and F1200 in comparison to composites containing 30 vol.%. Comparable values of approximately 2.1 J were observed for composites containing F80 particles.

Impact tests reveal a significant influence of particle size and content on the material’s ability to absorb energy. The main tendency visible in the data is an increase in impact strength as the size of the reinforcing particles decreases. The impact strength for the composite with 30 vol.% of F80 particles was 8.9 J/cm², while for the F150 fraction it increased to 22.9 J/cm², and for the finest F1200 particles it reached 36.5 J/cm². A similar trend was observed for composites with 40 vol.% particles. The increase in impact strength with decreasing particle size is consistent with the reinforcement mechanisms, which involve blocking and dispersing crack propagation by numerous fine particles.

The second meaningful relationship concerns the effect of the volume content of α−Al₂O₃ particles on impact strength. In most cases, increasing the ceramic content from 30 vol.% to 40 vol.% led to a decrease in impact strength. For example, for the F150 fraction, the impact strength decreased from 22.9 J/cm² to 20.9 J/cm², and for F1200 from 36.5 J/cm² to 31.6 J/cm².

The exception to this rule is the composite with the largest particles, F80, where the impact strength remained constant at about 8.7–8.9 J/cm² regardless of the volume content of the reinforcement.

SEM scanning of fracture surfaces and microstructural analysis of samples cut in a plane perpendicular to the fracture allowed for a detailed examination of crack initiation and propagation under dynamic loading. The cracks run in a manner characteristic of brittle fractures, with minimal signs of plastic deformation. These deformations were visible in the form of elongated fragments of the matrix material located between the particles, mainly in the tensile zone on the opposite side to the impact site.

Secondary cracks were also observed at the fractures, leading to delamination and fragmentation, running in a plane perpendicular to the main fracture (the crack paths are marked with lines)—Figure 19a–c. They were caused by the concentration of compressive stresses, mainly in the hammer impact zone. In these areas, on the surface of the fracture, there was also visible severe defragmentation of the material, numerous brittle cracks, as well as traces of local upsetting and spalling of particles from the matrix—Figure 19d. At the edge of the fracture, the presence of particles strongly embedded in the matrix was confirmed—Figure 19f. After the crack passed along the particle–matrix boundaries, adhesively bound matrix fragments were observed on the particle surfaces—Figure 19e. In contrast to the compression zone, on the opposite side, where tensile stresses occur, the fractures were smoother and flatter (Figure 19g).

In composites reinforced with F800 and F1200 particles, the crack’s tendency to propagate along the boundaries between micro-areas reinforced with particles and zones with increased matrix volume. The crack propagation paths and front are shown in Figure 20.

During impact, energy is not effectively dissipated in large plastic deformations but generates local stress fields around the particles. In the composites studied, the main fracture front branches out in an irregular pattern. In places where particle boundaries are a weak point, these stresses initiate additional micro- or macro-lateral cracks. When a particle is torn out of the matrix, a local void is created, which changes the stress distribution and causes the crack to branch.

4. Discussion

The results confirm that squeeze casting is an effective route for fabricating AlMg7 matrix composites with a high degree of densification, reaching relative densities above 98%. Reinforcement effects were evident in the hardness increase from 91 HBW for the base alloy to 194 HBW and in the rise in compressive strength σc_max, which reached 522 MPa for composites with 40 vol.% F1200 particles.

Composites reinforced with 30 vol.% F150 particles achieved an exceptionally low porosity of 0.64%, indicating the most favorable particle size relationship for effective infiltration. The general tendency is that the compressive strength ${\sigma }_{cmax}$ is higher than the tensile strength. With this particle size, ${\sigma }_{cmax}$ reached 406 MPa, and ${\sigma }_{max}$ was 195 MPa, which is typical for brittle-ductile materials. Increasing the volumetric content of ceramic particles from 30 vol.% to 40 vol.%, especially in combination with the F1200 fraction, led to a distinct increase in compressive strength. The increase in the compressive yield strength CYS of the composites results from the synergy of dispersion strengthening and grain boundary hardening described by the load transfer and Hall–Petch relationships, respectively. Composites with 40 vol.% reinforcement fractured on a shear plane at an angle of approximately 45°. The contribution of dislocation strengthening related to the difference in the coefficients of thermal expansion, ${\Delta \sigma }_{CTE}$, was not calculated separately due to the scope of this work. However, it has been included in the subsequent research program.

It should be clearly emphasized that for the measured CYS value of 161 MPa for the unreinforced AlMg7 matrix material, the strengthening resulting from grain refinement Δσ_H-P, determined using the Hall–Petch relationship, is 30.87 MPa. The remaining, dominant part of the matrix yield strength, amounting to approximately 130 MPa, is the result of intensive solid-solution strengthening. This high base value stems from the high magnesium content of 7.15 wt% Mg in the Al matrix, where dissolved atoms and intermetallic phase precipitates, e.g., the β phase, Mg₂Si, effectively impede dislocation movement, determining the fundamental strength of the AlMg7 alloy.

Strengthening results from Hall–Petch effects, as Al₂O₃ particles act as pinning sites during solidification, restricting grain boundary migration and refining the α-Al matrix.

SEM observations of fracture surfaces revealed a strong dependence of failure mode on particle size. In F80-reinforced composites, cracking occurred mainly along interfaces, with particle pull-out dominating, indicating a low-energy brittle–plastic fracture mode. In contrast, F1200 particles more frequently forced cracks to propagate through the AlMg7 matrix, a more energy-intensive path that increased impact strength. Nonetheless, impact resistance was significantly reduced by process-related defects such as brittle silica binder bridges and Mg₂Si compounds, acting as stress concentrators, as confirmed by EDS/EDX analysis. Porosity also played a critical role, rising to approximately 1.8% for 40 vol.% F80 composites and around 1.0% for F1200, acting as crack initiation sites. As a result, the total impact energy fell sharply, from 8.8 J for 30 vol.% F1200 to only 2.1 J for F80 composites at both 30 and 40 vol.% Al₂O₃. These results highlight that optimal strength properties require not only the appropriate particle size and content but also control over the quality of the preforms and the manufacturing process.

In conclusion, the composites produced exhibit a favorable combination of high hardness, compressive strength, and low density. However, their limited impact resistance highlights the need for further process optimization, with particular focus on improving the particle–matrix interface. Preliminary ballistic tests have shown encouraging trends, and results from ongoing numerical optimization will be reported in future studies.

5. Conclusions

The squeeze casting method ensured a high-quality microstructure in the AlMg7−Al₂O₃ composites by providing a homogeneous matrix characterized by a small grain size of approximately 17.7 μm. This contributed to the material’s strengthening, generating a Hall–Petch contribution of about 31 MPa in the pure matrix alloy. Rapid solidification under high pressure, combined with the blocking effect of the ceramic particles, prevented the formation of large, brittle β−Al3Mg2 phase precipitates. Ultimately, the process delivered low porosity of up to approximately 1% with a relative density above 98.9%. This confirms the high material integrity, except for composites with the F80 fraction, in which post-process defects were observed.

The strength of the composites is strongly dependent on the volume fraction and size of the particles. The highest values of static properties were achieved for 40 vol.% of F1200 particles, resulting in a hardness of 194 HBW, which represents a 113% increase relative to the pure matrix with a hardness of 91 HBW, and a compressive strength of 522 MPa. These results stem from synergistic structural strengthening, where load transfer (Δσ_L) contributes to strengthening of approximately 20 MPa for 30 vol.% and approximately 30 MPa for 40 vol.% of particles. The increase in strength is also consistent with the assumptions of the Orowan theory, which predicts greater strengthening effectiveness for composites with smaller (F1200) particles.

The highest compressive strength of 522 MPa was exhibited by composites with 40 vol.% F1200 particles, while the highest tensile strength of 341 MPa was shown by composites with 30 vol.%. This difference results from the fact that a lower ceramic phase content minimizes stress concentration and the risk of crack initiation, which is critical for properties under tensile loading. In the scope of dynamic properties, despite the general reduction in impact toughness of the composites relative to the matrix, the fine F1200 fraction significantly increases the energy absorption during fracture. The impact toughness, measured by the absorbed energy for 40 vol.% of F1200, was 8.8 J, which was approximately four times higher than for 40 vol.% of F80, which was 2.1 J.

The fracture behavior of the materials is highly dependent on the size and fraction of the particles, which regulates the overall brittleness and impact toughness. Large F80 particles promoted brittle interfacial fractures in both static and dynamic tests, as well as crack initiation at brittle Mg₂Si intermetallic phases. In contrast, fine F1200 particles led to more energy-intensive processes, such as fragmentation of the particles themselves and transcrystalline fractures in the matrix, which is a key factor explaining the higher energy absorption in composites with the fine fraction.